Method of fabricating nickel base superalloys having improved stress rupture properties

ABSTRACT

Improved superalloy powder materials consisting of a powder mixture of a substantially carbon free nickel base superalloy containing gamma prime phase and a metastable carbide such as vanadium carbide (VC), titanium carbide (TiC) or chromium carbide (Cr3C2) in an amount sufficient to provide a controlled carbon content in the resulting mixture. The mixture is then consolidated into a shape by heat and pressure, after which it may be mechanically worked at a temperature below the recrystallization temperature to thereby promote grain growth in the shape, followed by heat treating the shape at a temperature above the recrystallization temperature to solution the metastable carbide and permit grain growth. Finally, the shape is subjected to an aging treatment at a temperature below the recrystallization temperature to reprecipitate the gamma prime and carbide phases in the matrix and the new grain boundary areas.

United States Patent 1 Collins et al.

[ 1 Apr. 9, 1974 Lawrence D. Graham, Euclid; Charles S. Kortovich, Wickliffe, all

of Ohio [73] Assignee: TRW Inc., Cleveland, Ohio [22] Filed: Mar. 12, 1973 [21] Appl. No.: 340,177

[52] U.S. Cl 148/126, 29/4205, 75/200, 75/226. 148/115 F, 148/127 [51] Int. Cl B22f 3/16, C22c 19/00 [58] Field of Search 75/226, 214, 200; 148/126, 148/127. 162. 11.5 F; 29/4205 3.723.076 3/1973 Benesovsky 75/200 X Primary Examiner-Carl D. Quarforth Assistant E.\'aminerR. E. Schafer Attorney, Agent, or Firm-Hill. Sherman, Meroni,

Gross & Simpson 5 7 ABSTRACT Improved superalloy powder materials consisting of a powder mixture of a substantially carbon free nickel base superalloy containing gamma prime phase and a metastable carbide such as vanadium carbide (VC), titanium carbide (TiC) or chromium carbide (Cr C in an amount sufficient to provide a controlled carbon content in the resulting mixture. The mixture is then consolidated into a shape by heat and pressure. after which it may be mechanically worked at a temperature below the recrystallization temperature to thereby promote grain growth in the shape, followed by heat treating the shape at a temperature above the recrystallization temperature to solution the metastable carbide and permit grain growth. Finally, the shape is subjected to an aging-treatment at a temperature below the recrystallization temperature to reprecipitate the gamma prime and carbide phases in the matrix and the new grain boundary areas.

8 Claims, N0 Drawings METHOD OF FABRICATING NICKEL BASE SUPERALLOYS HAVING IMPROVED STRESS RUPTURE PROPERTIES The invention described herein was made in the performance of work under NASA Contract No. NAS 3l 3488 and is subject to the provisions of Section 305 of the National Aeronautics and Space Act of l958 (72 Stat. 435; 42 U.S.C. 2457).

BACKGROUND OF THE INVENTION 1. Field of the Invention This invention is in the field of fabricating superalloy materials with elevated temperature mechanical properties which are superior to those produced by other powder metallurgy techniques. It involves the use of a mixture of a substantially carbon free superalloy with a metastable carbide as a starting material, the powder mixture being formed into fully dense material utilizing normal consolidation procedures. The consolidated material is then given a thermal-mechanical treatment in order to achieve optimum grain size and microstructure. This treatment may consist first in adding mechanical energy to the material at temperatures below the recrystallization temperature, such as by forging in order to promote grain growth and produce a more uniform heat treatment response throughout the material. This is followed by heat treatment above the recrystallization temperature to solution the metastable carbides and grow the grains to the required size. The thermalmechanical treatment is concluded with a lower temperature aging treatment which produces the optimum microstructure by reprecipitating the gamma prime and carbide phases in the matrix and the new grain boundary areas.

2. Description of the Prior Art The metallurgy of nickel base alloys for high strength, high temperature applications has been developed for several decades and is largely based on the discovery that the elevated temperature mechanical properties of a solid solution nickel base alloy could be greatly enhanced by the precipitation of a finely dispersed second phase. Since that time, alloy development studies have produced vast improvements in the capabilities of nickel base alloys for service at higher temperatures, largely for the aircraft industry. In structural applications, involving such forms as sheet, forgings. and castings, material improvements can be translated into increased operating temperatures and reduced weight. In power plant operations, such as turbine blading for turbo jet engines, they can be translated into increased efficiency and power output.

One of the hardening mechanisms for nickel base superalloys is solid solution hardening which can be defined as an increase in resistance to deformation obtained by dissolving one element in another. This increase in resistance to deformation can be interpreted by means of dislocation theory. The movement of dislocations. or atomic disregistries, through a crystal lattice is the mechanism that permits plastic deformation in a metal or alloy. Any disturbance within the crystal lattice that hinders the movement of dislocations reduces the rate of plastic deformation, thus strengthening the alloy. The introduction of dissimilar atoms forces the crystal lattice of the solvent either to expand or to contract. This enforced change in lattice dimension causes a disturbance, which takes the form of a strained condition within the crystallographic structure of the alloy. Small areas of localized strain surround each of the solute atoms. It is the strained areas that retard the movement of dislocations and strengthen the alloy.

Chromium is generally added to nickel base superalloys intended for elevated temperature service. The addition of chromium serves to increase the strength by solid solution hardening, improves the oxidation resistance of the nickel, increases the recrystallization temperature range, and improves the sulfidation resistance of nickel.

Many modern day nickel base superalloys include not only chromium but also molybdenum, tantalum or tungsten as additional solid solution strengtheners.

The mechanism which is responsible for most of the high strength of nickel base superalloys at high temperatures is the precipitation hardening mechanism. The strengthening due to precipitation hardening depends upon the formation of a dispersed second phase within the solvent solution matrix. Two heat treating steps are usually necessary to obtain this structure in wrought nickel-base superalloys. The first step, termed solution treating or solution annealing involves heating a solid solution above the solvus temperature, holding for a suitable length of time, and then cooling to room temperature at a rate necessary to retain the elevated temperature structure. This treatment results in a supersaturated solution. The second step, termed aging, results in the formation of a dispersed second phase by precipitation from the supersaturated solid solution. Although the aging step can be carried out within a reasonable time at room temperature in some alloy systems, it is usually done at an intermediate temperature. The selected temperature represents a compromise between diffusion rate and degree of supersaturation. At room temperature, the diffusion rates in nickel base superalloys are so low that no precipitation can occur. Solute diffusion is a phenomenon vital to precipitation because clustering of solute atoms is a prerequisite to precipitation. The solute atoms must migrate, by diffusion, to concentrate or cluster at locally favorable sites and initiate the formation in the second phase. Diffusion rates increase with increasing. temperature. Thus, some intermediate temperature above room temperature favors the formation of the second phase. Increasing the temperature, however, usually increases the amount of solute that can be held in solid solution. Because the amount of second phase obtainable is inversely related to the amount of solute which can be held in solid solution, an increase in aging temperature decreases the amount of second phase. As a general rule, increased amounts of second phase result in increased strength properties. Thus, some intermediate temperature above room temperature decreases the amount of second phase. If a low temperature is selected, the result will be a large eventual volume of second phase, but it will be a long time before anything approaching equilibrium in the precipitation reaction is attained. If a very high temperature is selected, the result will be a small volume of second phase, but this volume will be attained in a short period of time.

In modern day nickel base superalloys, the second phase normally includes the intermetallic compound Ni Al, usually designated gamma prime located in a solid solution matrix, designated gamma. It is generally considered that up to about 65 percent of the aluminum of the intermetallic compounds can be replaced by titanium. Both the solid solution matrix and the Ni Al compound have a face centered cubic crystal structure. Although the lattice parameter of the matrix changes slightly with small changes in composition, the difference in lattice parameters between matrix and the second phase is very slight, generally not more than about 0.5 percent. However small this difference may be, it creates considerable internal strain during the period of coherency when the lattice of the precipitate is forced to conform to the lattice of the matrix. The gamma prime phase is capable of taking a considerable amount of other elements into solution such as chromium, titanium and cobalt. The controlled precipitation of gamma prime in the matrix is thus of critical importance in proper strengthening of nickel base superalloys.

As the strength levels of nickel base superalloy systems were increased, the workability of wrought alloys became severely limited, resulting in the necessity of using investment casting processes for the production of complex shapes. Alloy development programs utilizing investment casting techniques have resulted in a complex series of alloys being produced. As these alloys have increased in complexity, i.e., through larger amounts of more alloying additions, their castability has become a serious problem and the degree of segregation has tended to increase. Alloy segregation has led to undesirable scatter in the mechanical properties.

Powder metallurgy techniques offer the potential of fabricating superalloy materials with elevated temperature mechanical properties equivalent to or even superior to currently available cast or wrought materials. At the present time, however, this potential has not been achieved at temperatures above about 1,400F (760C). This is due primarily to the inability to attain the grain size and microstructure required to develop the high temperature properties. This inability appears to be due, at least in part, to the presence of stable carbides which form either during the melting process or during the consolidation process. Such carbides can act as pinning points thereby retarding grain growth and in general do not respond satisfactorily to heat treatments, thereby making it difficult to obtain the optimum microstructure.

SUMMARY OF THE INVENTION In general, the present invention involves mechanically mixing metastable carbide powders such as vanadium carbide (VC), titanium carbide (TiC) or chromium carbide (Cr C with a nickel base superalloy powder which is essentially free from carbon or at least has a carbon content not in excess of about 0.1 percent by weight in order to bring the carbon concentration to a level which is dictated by the composition of the base material. Generally, this level does not exceed about 0.5 percent by weight and is preferably from 0.2 to 0.3 percent by weight. The powder mixture is formed into fully dense material utilizing normal consolidation procedures such as extrusion or hot isostatic pressing. The consolidated material is then given a thermalmechanical treatment in order to produce an optimum grain size and microstructure. The treatment consists first of adding mechanical energy to the material at temperatures below the recrystallization temperature, preferably from 50 to 150F (-66C) below, such as by a forging operation, in order to promote grain growth and produce a more uniform heat treatment response throughout the material. Up to a 10 percent reduction in diameter may be achieved in this working. Next, a heat treatment above the recyrstallization temperature is employed to solution the metastable carbides and grow the grains to the required size. The preferred solution temperature is from 2250 to 2425F (l,232l,330C). The thermal-mechanical treatment is concluded with a lower temperature aging treatment (typically l,500 to 2,000F (816-l,093C)) which secures the optimum microstructure by reprecipitating the gamma prime and carbide phases in the matrix and new grain boundary areas. The amount of mechanical energy required in the heat treatment temperatures depends upon the base composition, the amount of carbide added, and the consolidation process.

For the purposes of this invention a metastable carbide is defined as one having a free energy of formation (AG) at l300K of from 11 to 41 kilocalories per mole.

Another feature of the present invention is the addition of small amounts of hafnium or tantalum to those nickel base superalloys which do not already contain these elements. The addition of carbon in the form of metastable carbides results in their decomposition during grain growth heat treatments without impairing the growth of the grains themselves. The carbon thus released is available for subsequent precipitation of MC type hafnium and tantalum rich carbides to enhance strength and ductility.

DESCRIPTION OF THE PREFERRED EMBODIMENTS The process aspects of the present invention are applicable to nickel base superalloys generally in which a gamma prime phase is capable of precipitating into a matrix of gamma phase. Such nickel base superalloys usually have compositions falling within the following relatively broad analyses:

Cr 5.0 25 Co 5.0 25 Mo to l2 W to l5 Cb to 4 Ti 0.5 6 Al 1 7 B to 0.4 Zr to 0.3 Ta to 12 Hf to 5 Re to 1 Within the broader ranges, the following ranges of ingredients are particularly preferred:

Cr 6 20 Co 6 13 Mo to ID W to 13 Cb to 2.0 Ti 1 5 Al [.3 5.9 B 0.005 0.2 Zr 0.05 0.2 Ta 1 9 Hf 0.3 4 Re to 0.5

In connection with the preferred ranges, it is particularly preferred that the sum of the molybdenum, tungsten and tantalum contents range up to 17 percent, and that the sum of the aluminum and titanium contents be within the range of about 6 to 12 percent.

There are numerous nickel base superalloys presently available commercially to which the present application is applicable. and the nominal compositions of such superalloys. with the carbon content that they with a break-through pressure of approximately 100,000 psi (689.5 MN/m).

All four alloys were successfully extruded and radiographic inspection indicated no visible defects. Metalnormally contain, is given 1n the following tabl 5 lographic analyses of the as-extruded bars revealed AllOY ('1 M0 w Ch '11 Al H Zr ()thcr MAR-M2011 0 15 0.0 10.0 12.5 1.0 2.0 5.0 0.015 0.05

B-l900 l1 1H 8.0 10.0 6.0 0.6 5.9 0.01 0.10 4.3 Ta 00100 015 10.0 15.0 3.0 4.7 5.5 0.015 0.00 1.0 v MAR-M246 0 15 0.0 10.0 2.5 10.0 1.5 5.5 0.015 0.05 1.5 Ta Rcl'lt: 4| 0 09 19.0 11.0 10.0 3.1 1.5 0.01

RCHL 05 0.15 14.0 x0 3.5 3.5 3.5 2.5 3.5 0.01 0 05 l'dlmcl 500 0.024 10.0 111.0 4.0 2.9 2.9 0.005

lthmct 700 0.15 15.0 18.0 4.0 3.5 4.3 0.030

Waspniux 0.03 195 I35 4.3 3.0 l.3 0.006 0.06

Amount 100 Percent- Carbon Mush Base Amount age Carbon Content lzxlru Alluy (nrhidc in Carbide of Alloy 0011 N0 t m; (gms) (w/o) (vi/ 1 1 mm .7 V 0 2 75 7 54.3Cr, 13.2 .20 1 1592 x 36. TiC 19.5 .20 4 .1502 38VC 18.3 .20

The alloy constituents were placed in a welding chamber along with a double conical blender. The welding chamber was then back-filled with argon and the proper amount of 6 to 8 micron Cr c 3 to 6 micron TiC, or 3 to 4 micron VC was sealed in the conical blender along with the base powder. The blender was then removed from the welding chamber, mounted on a lathe and rotated for one hour before transfer into the welding chamber along with a stainless steel extrusion can.

After the welding chamber was again back-filled with argon. the powder was removed from the blender and placed into the extrusion can (approximately 75 percent dense) and covered with a 16 gauge stainless baffle. After pumping down overnight to a vaccum of approximately 0.2 microns, the chamber was back-tilled with argon and the nose of the extrusion can welded into place. The extrusion can was then transferred to an electron beam welder where a hole was made in the nose and the chamber pumped down overnight to a vacuum of about 0.1 micron or less. As the gas escaped during pumpdown, the baffle prevented powder from being pulled out through the hole. After evacuation the electron beam was defocused and the hole rescaled.

The extrusion cans were then preheated to l,500F (815C) for one hour in a resistance element heated box furnace and transferred to a barium chloride salt bath at 2.175F (l.l90C) and soaked for one hour prior to extrusion. Single extrusion was conducted in a 700 ton hydraulic extrusion press using zirconia coated tool steel dies at a 16 to 1 reduction ratio. The ram speed was approximately 2 inches (5.1cm) per second considerable differences among the four extrusions. The as-extruded base material evidenced recrystallization only in central portions of the bar stock and had only a limited amount of banding, primarily at the outer portions. The central portion of the bar had a grain size of ASTM No. 8 (25 microns), and while limited amounts of a gamma prime type precipitation was observed throughout the matrix, a considerable amount was seen in the grain boundary areas.

No recrystallization was observed in the material containing Cr C A considerable amount of banding was noted in the microstructure and was attributed to the large amounts of small particles (about 2 microns) seen throughout the matrix. Numerous large (about 24 microns) particles were also seen which appeared to be similar to those smaller precipitates responsible for banding. Microprobe analysis of these larger particles indicated that they were MC type carbides.

Recrystallization to an ASTM No. 7 (35 microns) grain size occurred in the alloy containing TiC, but only in the central portion of the extrusion. There was a moderate amount of banding throughout the microstructure, but the large particles found in the extrusion containing chromium carbide were not observed.

Recrystallization to an ASTM No. 7 grain size was observed in the material containing VC but only in the central portion of the extrusion. There was a considerable amount of handing throughout the microstructure and the original carbides also appeared somewhat agglomerated.

The tensile and stress rupture properties for the four alloys in the as-extruded condition as well as those for the comparable cast alloy were obtained. The ultimate tensile strength and the 0.2 percent offset yield strength of all four powder alloys at room temperature and at 1,400F. (760C.) exceeded that of the cast alloy. This can be attributed to the extremely fine grain size (ASTM 7-8) of the extruded powder. The percent elongation for all four powder alloys at room temperature exceeded that of the cast alloy. Values were slightly lower at 1,400F (760C) with the exception of material containing Cr C which was slightly higher.

In order to increase the grain size in the extruded alloys. but avoid excessive amounts of porosity, a series of vacuum heat treatments were developed to promote diffusion of the internal gases out of the microstructure. These treatments included multiple exposures starting below incipient melting to produce the proper grain size. As a result of these vacuum heat treatments, a grain growth treatment was selected for each alloy, and this is given in the following table:

2 hours at Z400F l3 lbC) vacuum. 4 5 step aging (l) Fiver step aging: 4 hours at 2000"l- (l093(')/h hours at I600F (870(')/4 hours at 1X00F (982(l/24 hours at |200F (64l()!8 hours at I400F (760%) An additional series of extrusions were made up from powder material which had been atomized in hydrogen. Two different powder sizes were used, one ranging from 60 to +100 mesh, the other from -l to +325 mesh. The constituents for these extrusions are given below:

content of 0.2 weight percent, the processing consisted of the side press with a 24 hour heat treatment at 2,355F (l,290C) in argon. For the alloy containing vanadium carbide in an amount sufficient to provide an overall carbon content of 0.15 weight percent, the treating schedule consisted of side pressing followed by Amount Amount 100+ 125 -fi0+l00 Amount Percentage Percentage Intru ion Bust: Alloy Base Allo Carbide Carbon in Carbon in :No (gms) (gms) (gms) (arbitlc Alloy h 2270 I324 37 ZTlC I95 0.20 7 2270 l342.4 lQliVC [3.3 ll.| 8 2270 [3]}.4 49.6VC [3.3 0. 8

The extrusion procedure was essentially similar to that previously described except that nitrogen gas was used to backfill the chamber instead of argon in the first instance. The extrusions were made at 2,175F. l,l90C) at a 16 to 1 reduction ratio with the following extrusion parameters:

mechanical working as a means of promoting a more uniform heat treatment response. Side pressing forging runs were conducted on 9% inch (1.27 cm) long specimens of each alloy. The specimens were first double coated with a lubricant to offer protection during heatup to the forging temperature. Side press forging was conducted on a 150 ton hydraulic forging press at temperatures below the recrystallization point for each alloy and deformations of approximately 5 percent reduction in diameter were imparted in one blow.

The controlled deformation was followed by high temperature heat treatments for grain growth. As a result of further studies, a heat treatment was selected for each of the alloys (extrusions Nos. 6 to 8) which produced the desired ASTM No. 2 to 3 grain size in the microstructure. For the alloy containing titanium carbide in an amount sufficient to produce an overall carbon 24 hour heat treatment at 2,345F (1,285C) in vacuum, followed by a 24 hour heat treatment at 2,355F (1,290C) in vacuum. For the alloy containing vanadium carbide in an amount sufficient to produce an overall carbon content of 0.28 weight percent, the treating schedule consisted of side pressing and a 24 hour heat treatment at 2,400F (1,316C) in vacuum.

As a result of these studies, we have concluded that improved properties were achieved by using a base alloy (carbon free Mar-M246) with enough vanadium carbide to provide a carbon content of 0.28 weight percent. Powder consolidation to eliminate incomplete compaction was conducted in a two step operation consisting of hot pressing followed by hot extrusion. Prior to hot pressing, the powders were weighed out and blended in a nitrogen atmosphere for two hours, canned under a protective covering of argon, vacuum washed at l,OO0F (540C) for 24 hours and then forged at 2,l50F (1,175C) to achieve a 5 percent reduction in diameter. Following this, the material was given various solution heat treatments. A typical heat treatment consisted of 24 hours at 2,380F (1,305C) in argon plus 24 hours at 2,390F (l,3lOC) in argon prior to homogenization at 2,400F (l,316C) thereby resulting in the desired ASTM No. 2-3 grain size. The solution heat treatment above the recrystallization temperature of the alloy was then followed by an aging treatment below the recrystallization temperature, specifically by a 1,975F (925C) age for 4 hours, and a l,700F (925C) age for hours. The mechanical properties resulting from alloys so treated are given in the following table:

Tensile Properties Stress Rupture Properties Temperature Ultimate Strength YiclcTstrcngth Elonf": R.A."/1 Loud Llfe Elan/71 R.A.%

F (ksi) (ksi) (ksi) (hrs.)

Room 124.3 110.0 2.9 3.7 1400 135.3 120.6 2.1 4.3 114.5 110.7 1.3 1.6 1900 40.2 33.5 5.3 6.6

Cast MAR-M24o Room 140 125 5.0

1') Estimated l-mm lursun-Miller Plot The alloy described above represents, we believe, the best combination offered by conventional powder metallurgy processing to date.

The alloys of the present invention can also be improved by adding carbide formers such as hafnium and tantalum in amounts up to about 1 percent where the superalloy composition does not include these carbide formers. It is also possible to add additional gamma prime formers and solid solution hardeners to the alloy without the segregation problems experienced with casting alloys.

The alloy produced according to the present invention has an equiaxed grain shape of approximately ASTM No. l-3 grain size and a discrete particle type carbide and gamma prime morphology at the grain boundaries.

It should be evident that various modifications can be made to the described embodiments without departing from the scope of the present invention.

We claim as our invention:

1. The method of fabricating a nickel base superalloy containing gamma prime phase which comprises mixing a powder of said superalloy substantially devoid of carbon with a metastable carbide powder in proportions sufficient to achieve a significant carbon content in the resulting mixture, which carbon content is not in excess of about 0.5% by weight, consolidating the resulting mixture into a shape by heat and pressure, mechanically working the resulting shape at a temperature below the recrystallization temperature as required to promote grain growth in the shape, and thereafter heat treating the shape at a temperature above the recrystallization temperature to solution the metastable carbide and promote grain growth.

2. The method of claim 1 in which the carbon content of said mixture is in the range from about 0.2 to 0.3 percent by weight.

3. The method of claim 1 in which the carbon content of the superalloy initially is not in excess of about 0.1 percent by weight.

4. The method of claim 1 in which the grain size of the product after heat treating is such that the majority of the grains are in the range of 2 to 3 on the ASTM scale.

5. The method of claim 1 in which said heat treatment is followed by an aging treatment at a temperature of below the recrystallization to repreeipitate the gamma prime and carbide phases in the matrix and new grain boundary areas.

6. The method of claim 1 in which said superalloy has a composition within the following ranges:

Cr 5.0 25 Co 5.0 25 Mo to 12 W to 15 Cb to 4 Ti 05 6 A1 1 7 B to 0.4 Zr to 0.3 Ta to 12 Hf to 5 Re to l 7. The method of claim 1 in which said superalloy ha a composition within the following ranges:

Cr 6 20 Co 6 13 Mo to 10 W to 13 Cb to 2.0 Ti 1 5 A1 1.3 5.9 13 0.005 0.2 Zr 0.05 0.2

Ta 1 9 Hf 0.3 4 Re to 0.5 

2. The method of claim 1 in which the carbon content of said mixture is in the range from about 0.2 to 0.3 percent by weight.
 3. The method of claim 1 in which the carbon content of the superalloy initially is not in excess of about 0.1 percent by weight.
 4. The method of claim 1 in which the grain size of the product after heat treating is such that the majority of the grains are in the range of 2 to 3 on the ASTM scale.
 5. The method of claim 1 in which said heat treatment is followed by an aging treatment at a temperature of below the recrystallization to reprecipitate the gamma prime and carbide phases in the matrix and new grain boundary areas.
 6. The method of claim 1 in which said superalloy has a composition within the following ranges:
 7. The method of claim 1 in which said superalloy has a composition within the following ranges:
 8. In a method of producing a nickel base superalloy from a powder material by compaction and heat treatment, the improvement which comprises using as said powder material a powder mixture of a substantially carbon free nickel base superalloy containing gamma prime phase and a metastable carbide having a free energy of formation, Delta G, at 1,300*K in the range from 11 to 41 kilocalories per mole in an amount sufficient to provide a carbon content not in excess of about 0.5 percent in the resulting mixture. 